Steel materials and a method for manufacturing the same

Abstract

Provided are: a steel material that contributes to the realization of both high strength and hydrogen embrittlement resistance; and a method for producing the steel material. This steel material has a chemical composition of: 0.15 to 0.35% of C; 0.8 to 2.5% of Si; 0.8 to 2.5% of Mn; 0.03 to 2.0% of Al; 0.002 to 0.010% of N; 0.01% or less of P; 0.01% or less of S; 0.01% or less of O; 0.0001 to 0.005% of B; 0.0 to 0.05% of Nb; 0.0 to 0.2% of Ti; 0.0 to 0.05% of V; 0.0 to 1.0% of Mo; 0.0 to 1.0% of Cr; 0.01 to 1.0% of Ni; 0.05 to 1.0% of Cu; 0.0005 to 0.01% of at least one of Ca, Mg, and REM; and the balance being Fe and impurities. The steel material has a martensite phase or a bainite phase containing dispersed and precipitated carbides.

Description

Technical Field

[0001] The present invention relates to steel materials used in all industrial fields such as automobiles, building materials, machine parts, home appliances, hydrogen stations, high-strength bolts, etc., and a method for manufacturing the same.

Background Art

[0002] (Description of related applications) The present invention is based on a claim of priority from Japanese Patent Application: Japanese Patent Application No. 2020-57734 (filed on March 27, 2020), and the entire description of the application is incorporated herein by reference. From the viewpoints of weight reduction of vehicle body parts, reduction of construction costs, and response to global environmental problems, further strengthening of steel materials (steel sheets) and improvement of hydrogen embrittlement resistance characteristics are required. In addition, in fields such as industrial machinery, tanks, and line pipes, as the strength of steel materials increases, the operating environment becomes more severe. It is known that the strengthening of steel materials and the severity of the operating environment increase the hydrogen embrittlement susceptibility (HE susceptibility) of steel materials, and the development of high-strength steel materials with excellent hydrogen embrittlement resistance characteristics is required.

[0003] Under such circumstances, Patent Documents 1 and 2 disclose technologies for improving the high strength and hydrogen embrittlement resistance characteristics of steel sheets and steel materials by optimizing components, controlling carbide precipitation, and optimizing heat treatment. Further, Patent Document 3 discloses the world's first technology for strengthening and toughening low-alloy steel by dispersive precipitation of ε-carbide. Additionally, Patent Document 4 discloses a steel sheet that is excellent in all of tensile strength, ductility, hole expansion property, hydrogen embrittlement resistance property, and toughness by means of optimizing components, refining crystal grains, controlling the steel structure, and hydrogen trapping by finely dispersed carbonitrides. Furthermore, Patent Document 5 discloses an idea for high-strength bolts in which La, which is a kind of REM (Rare Earth Metal), is added to steel to precipitate LaNi5 in the steel, trap the hydrogen invading from the outside inside the crystal of La5Ni, and improve the hydrogen stress corrosion cracking resistance.

Prior Art Documents

Patent Documents

[0004]

Patent Document 1

Patent Document 2

Patent Document 3

Patent Document 4

Patent Document 5

Non-Patent Documents

[0005]

Non-Patent Document 1

Non-Patent Document 2

Non-Patent Document 3

Non-Patent Document 4

Non-Patent Document 5

Non-Patent Document 6

Non-Patent Document 7

Non-Patent Document 8

Summary of the Invention

Problems to be Solved by the Invention

[0006] The following analysis is provided by the inventor of the present application.

[0007] However, with the technologies described in Patent Documents 1 to 5, it has been difficult to achieve both high strength and hydrogen embrittlement resistance characteristics, or the achievement has been insufficient. For example, high-strength steel sheets with a tensile strength of 1180 MPa or more are applied around the cabin of a passenger car, but there is concern about delayed fracture during use. Also, regarding high-strength bolts for automobiles, due to the risk of fracture caused by hydrogen embrittlement, the adoption of high-strength bolts with a strength of 1200 MPa or less has been limited. Furthermore, in Patent Document 3, not only is the cost increased because expensive additive elements are required, but also the perspective on hydrogen embrittlement is lacking.

[0008] The main problem of the present invention is to provide a steel material and a method for manufacturing the same that can contribute to achieving both high strength and hydrogen embrittlement resistance characteristics.

Means for Solving the Problems

[0009] The steel material according to the first aspect has a chemical composition represented by, in mass %, C: 0.15% to 0.35%, Si: 0.8% to 2.5%, Mn: 0.8% to 2.5%, Al: 0.03% to 2.0%, N: 0.002% to 0.010%, P: 0.01% or less, S: 0.01% or less, O: 0.01% or less, B: 0.0001% to 0.005%, Nb: 0.0% to 0.05%, Ti: 0.0% to 0.2%, V: 0.0% to 0.05%, Mo: 0.0% to 1.0%, Cr: 0.0% to 1.0%, Ni: 0.01% to 1.0%, Cu: 0.05% to 1.0%, at least one of Ca, Mg and REM: 0.0005% to 0.01%, and the balance: Fe and impurities, and has ε-carbide with a size of 2 nm or more and 150 nm or less dispersed and precipitated at a density of 1×10 2 or more per mm 6 in the martensite phase or bainite phase having the above density.

[0010] The manufacturing method of the steel material according to the second aspect is a manufacturing method of the steel material for manufacturing the steel material according to the first aspect. When cooling from after hot rolling the steel material to room temperature, cooling is performed so that the martensite phase or bainite phase is formed, thereby dispersing and precipitating aluminum nitride in the martensite phase or bainite phase, and then tempering is performed at 100°C or more and less than 300°C to grow the ε-carbide with the aluminum nitride as nuclei. The method includes the steps of:

Effect of the Invention

[0011] According to the first and second aspects, it is possible to contribute to achieving both high strength and hydrogen embrittlement resistance characteristics.

Brief Description of the Drawings

[0012]

Figure 1

Embodiment for Carrying Out the Invention

[0013] In order to solve the above problems, the inventors of the present application studied the precipitation process of ε-carbide in the iron-carbon alloy system from an atomic perspective. As a result, they focused on the fact that when the martensite phase and / or bainite phase is formed (corresponding to the transformation from the austenite phase to the martensite phase or bainite phase), the solid solution amount of nitrogen atoms decreases to 10% or less (Non-Patent Document 5). By utilizing this, it was found that a structure in which aluminum nitride phases are finely dispersed and precipitated can be formed, and at the subsequent tempering treatment stage, ε-carbide can be dispersed and precipitated with these aluminum nitrides as nuclei. Note that both the ε-carbide and the aluminum nitride phase belong to the hexagonal crystal system, and the lattice constants of both are in a relationship where the formation of a semi-coherent interface is possible. Further, as a result of conducting various tests on steel materials having a structure in which ε-carbide is dispersed, the inventors of the present application found that high strength and extremely excellent hydrogen embrittlement resistance can be achieved simultaneously. Based on these findings, the inventors of the present application further conducted intensive studies and as a result, arrived at the present disclosure shown below.

[0014] Here, the inventors of the present application have studied the basic physical properties of ε-carbide in Fe-C alloys, which have been overlooked until now because they are transient carbides, using the method of frequency sweep type mechanical spectroscopy and published a paper (Non-Patent Document 1). ε-Carbide in steel has a lattice structure containing many carbon atom vacancy sites and has something in common with known substances known as hydrogen storage materials. Considering a report (Non-Patent Document 2) that an experiment was conducted to test the hypothesis that ε-carbide absorbs hydrogen up to the composition of Fe₂CH in the passivation study of steel, it is inferred that hydrogen will occupy the carbon atom vacancy sites of ε-carbide. It was pointed out in Non-Patent Document 3 that ε-carbide dispersed and precipitated in the ferrite phase can improve the strength of the ferrite phase by the particle dispersion mechanism. Also, Non-Patent Document 4 showed that ε-carbide precipitates with TiC as the nucleation site in rapidly cooled Fe-C-Ti alloys. In the prior art, ε-carbide has only been described additively at the end of a group of iron carbides as a superordinate concept, and its function and actual precipitation conditions have been ignored. Moreover, the idea of using ε-carbide as a substance that absorbs hydrogen invading steel has never existed. Therefore, steel materials that attempt to achieve both high strength and hydrogen embrittlement resistance by actively dispersing and precipitating ε-carbide in steel materials with a martensite phase or bainite phase as the main structure have not been developed.

[0015] In the present disclosure described below, the steel materials according to the following Form 1 and their modified forms can be appropriately selected and combined.

[0016] As the steel material according to the above Form 1, by mass, C: 0.15% to 0.35%, Si: 0.8% to 2.5%, Mn: 0.8% to 2.5%, Al: 0.03% to 2.0%, N: 0.002% to 0.010%, P: 0.01% or less, S: 0.01% or less, O: 0.01% or less, B: 0.0001%~0.005%, Nb: 0.0% to 0.05%, Ti: 0.0% to 0.2%, V: 0.0%~0.05%, Mo: 0.0% to 1.0% Cr: 0.0% to 1.0%, Ni: 0.01% to 1.0% Cu: 0.05% to 1.0% At least one of Ca, Mg, and REM: 0.0005% to 0.01% and, The balance: Fe and impurities and having a chemical composition represented by ε carbides with a size of 2 nm to 150 nm are 1 mm 2 1 x 10 per 6 The martensite phase or bainite phase is dispersed and precipitated at a density of 1000 or more. In the following, unless otherwise specified, mass% will be simply expressed as %.

[0017] Here, C (carbon) is an essential element that not only enables the phase transformation of steel but also improves the strength characteristics and hydrogen embrittlement resistance of steel by precipitating ε carbide. To obtain such effects, the C content must be 0.15% or more, preferably 0.17% or more, and more preferably 0.2% or more. On the other hand, if the C content exceeds 0.35%, the toughness and weldability of the steel will be significantly reduced, so the C content must be 0.35% or less, preferably 0.32% or less, and more preferably 0.3% or less. Note that ε carbide has a hexagonal crystal structure, and its composition is described in a steel materials textbook as Fe 2.4 C, but here we take into account non-stoichiometric compositions and use Fe 2.4~2.7 Let's call it C.

[0018] Si (silicon) not only dissolves in steel and contributes to the improvement of the strength of steel, but also has the effect of expanding the stable existence range of ε-carbide to the high-temperature side. To obtain such an effect, the Si content needs to be 0.8% or more, preferably 1.0% or more, and more preferably 1.2% or more. On the other hand, when the Si content exceeds 2.5%, the rolling load increases, so the Si content needs to be 2.5% or less, preferably 2.3% or less, and more preferably 2.0% or less.

[0019] Mn (manganese) dissolves in steel to strengthen it and suppresses the transformation to the ferrite phase during cooling, thereby lowering the Ms point (the temperature at which the transformation to the martensite phase starts during quenching). To obtain such an effect, the Mn content needs to be 0.8% or more, preferably 1.0% or more, and more preferably 1.2% or more. On the other hand, when the Mn content exceeds 2.5%, the weldability deteriorates, so the Mn content needs to be 2.5% or less, preferably 2.3% or less, and more preferably 2.0% or less.

[0020] Al (aluminum) is a useful element used as a deoxidizer during steelmaking. Al dissolved in steel combines with the dissolved nitrogen atoms during or after the transformation from the austenite phase to the martensite phase, bainite phase, or ferrite phase, and precipitates finely as aluminum nitride (AlN) inside the laths or along the dislocation lines, and its shape is plate-like or rod-like (Non-Patent Document 6). In the tempering treatment of steel, the aluminum nitride finely precipitated in the martensite phase or bainite phase functions as a nucleation site for ε-carbide. To fix the dissolved N as aluminum nitride, the Al content should be 0.03% or more, preferably 0.04% or more, and more preferably 0.05% or more. Also, considering the difference in the diffusion mobility of Al atoms and N atoms in steel to finely disperse and precipitate aluminum nitride, it is desirable that the ratio of "Al / N" in mass% of Al and N is greater than 7. On the other hand, when the Al content exceeds 2.0%, the inclusions in the steel increase and the ductility of the steel material decreases. Therefore, the Al content needs to be 2.0% or less, preferably 1.8% or less, and more preferably 1.5% or less.

[0021] N (nitrogen atom) is an essential element for forming aluminum nitride. However, when it exists in solid solution in steel, it is an element that reduces the toughness of the base material, so it is preferably reduced. It is known that the amount of N in solid solution decreases by one digit during the transformation from the austenite phase to the martensite phase, bainite phase, or ferrite phase (Non-Patent Document 5). Also, Non-Patent Document 6 reports that in the martensite structure, N reacts with Al to form aluminum nitride finely and uniformly. From the perspective of manufacturing cost, the N content can be 0.002% or more, but it may also be 0.003% or more, and further 0.004% or more. On the other hand, from the perspective of preventing strength variations, the N content should be 0.010% or less, preferably 0.008% or less, and more preferably 0.006% or less.

[0022] P (phosphorus) is an element that segregates at grain boundaries, weakens the grain boundary strength, and deteriorates the stress corrosion cracking resistance. Therefore, it is desirable to reduce the content of P as much as possible, but up to 0.01% is acceptable, preferably 0.005% or less, and more preferably 0.001% or less.

[0023] S (sulfur) is an element that forms MnS in steel and tends to be a starting point for stress corrosion cracking. Therefore, it is desirable to reduce the content of S as much as possible, but up to 0.01% is acceptable, preferably 0.005% or less, and more preferably 0.001% or less.

[0024] O (oxygen atom) is an element that combines with other elements to form oxides, leading to a decrease in the formability and toughness of steel. Therefore, it is desirable to reduce the content of O as much as possible, but up to 0.01% is acceptable, preferably 0.005% or less, and more preferably 0.001% or less.

[0025] B (boron) is an element that segregates at grain boundaries, increases the grain boundary strength, improves the toughness and stress corrosion cracking resistance, and significantly contributes to improving hardenability. To obtain such effects, it is desirable that the content of B contains 0.0001% or more, preferably 0.0005% or more, and more preferably 0.001% or more. On the other hand, when the content of B exceeds 0.005%, it precipitates as boride and deteriorates the toughness. Therefore, the content of B needs to be 0.005% or less, preferably 0.003% or less, and more preferably 0.002% or less.

[0026] Nb (niobium), Ti (titanium), and V (vanadium) are all elements that precipitate as carbonitrides during the tempering of steel, stably exist up to higher temperatures than ε-carbide, and contribute to maintaining strength. Nb, Ti, and V do not necessarily have to be contained, but at least one of Nb, Ti, and V can be selected as needed. The content of each of Nb, Ti, and V can be 0.0% or more, but in order to obtain the effect of maintaining strength, the content of each of Nb, Ti, and V is preferably 0.001% or more, more preferably 0.005% or more. On the other hand, when the content of each of Nb and V exceeds 0.05%, the toughness of the welded part decreases and the raw material cost increases, so the content of each of Nb and V needs to be 0.05% or less, preferably 0.04% or less, and more preferably 0.03% or less. The content of Ti can be tolerated up to 0.2%, but from the perspective of weldability, it is preferably 0.18% or less, and more preferably 0.15% or less.

[0027] Mo (molybdenum) and Cr (chromium) are both elements that dissolve in steel and contribute to improving the strength of the steel, or suppress the transformation to the ferrite phase during cooling to improve hardenability. Also, Mo and Cr form complex carbonitrides during the formation of the respective carbonitrides of Nb, Ti, and V. Mo and Cr do not necessarily have to be contained, but at least one of Mo and Cr can be selected as needed. The content of each of Mo and Cr can be 0.0% or more, but in order to obtain the effect of improving strength, the content of each of Mo and Cr is preferably 0.01% or more, more preferably 0.02% or more. On the other hand, when the content of each of Mo and Cr exceeds 1.0%, the hot workability of the base metal decreases, so the content of each of Mo and Cr needs to be 1.0% or less, preferably 0.8% or less, and from the perspective of cost, it is preferably 0.5% or less.

[0028] Ni (Nickel) is an austenitizing stabilizing element, which has the effect of suppressing hydrogen intrusion and is also effective in improving the stress corrosion cracking resistance. To obtain these effects, the Ni content needs to be 0.01% or more, preferably 0.02% or more, and more preferably 0.05% or more. On the other hand, when the Ni content exceeds 1.0%, not only do these effects saturate, but the cost also increases. Therefore, the Ni content is desirably 1.0% or less, preferably 0.8% or less, and more preferably 0.5% or less.

[0029] Cu (Copper) is an element that has the effect of suppressing hydrogen intrusion into steel materials and improving the stress corrosion cracking resistance. To obtain such an effect, the Cu content needs to be 0.05% or more, preferably 0.08% or more, and more preferably 0.1% or more. On the other hand, when the Cu content exceeds 1.0%, the hot workability of the base material decreases, so the Cu content needs to be 1.0% or less, preferably 0.8% or less, and more preferably 0.5% or less.

[0030] Ca (Calcium), Mg (Magnesium) and REM (Rare Earth Metal: Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb, Lu) all have a stronger affinity for S than Mn and can each form Ca-based sulfides, Mg-based sulfides or REM-based sulfides in steel, effectively contributing to the reduction of MnS, which is likely to be the starting point of stress corrosion cracking. For Ca, Mg and REM, at least one of Ca, Mg and REM can be selected. To obtain such an effect, the total content of Ca, Mg and REM needs to be 0.0005% or more, preferably not less than 0.001%, and more preferably not less than 0.002%. On the other hand, when the total content of Ca, Mg and REM exceeds 0.01%, the cleanliness of the steel decreases. Therefore, the total content of Ca, Mg and REM needs to be 0.01% or less, preferably 0.008% or less, and more preferably 0.005% or less.

[0031] The balance is Fe and impurities (inevitable impurities). However, as long as the effects of the present invention are not impaired, it does not exclude the inclusion of components other than the above.

[0032] Here, if the cooling rate is appropriately selected, ε-carbide precipitates during or after the transformation into the martensite phase or bainite phase, with the aluminum nitride precipitated earlier serving as the nucleation site in the tempering process. As a result, hexagonal ε-carbide with a size of 2 nm or more and 150 nm or less is dispersed and precipitated at a density of 1×10 2 or more per mm 6 in the martensite phase or bainite phase, obtaining a steel material. Since the previously precipitated aluminum nitride serves as the precipitation site, the ε-carbide can be controlled to the above density. The size can be controlled by adjusting the tempering temperature and time. The finely dispersed and precipitated ε-carbide contributes to the improvement of strength and ductility by the particle dispersion mechanism disclosed in Non-Patent Document 3. From the perspective of hydrogen embrittlement resistance, since the ε-carbide has a crystal structure capable of absorbing a large number of hydrogens (Non-Patent Documents 1 and 2), the hydrogen absorption amount is significantly larger than the amount of hydrogen captured at the interface between the conventional matrix structure and carbonitride, and the hydrogen embrittlement resistance of the steel material is dramatically improved. At this time, whether the interface between the ε-carbide and the matrix crystal is coherent or incoherent does not have a significant impact. Note that a part of Fe in the ε-carbide may be substituted with other elements such as Si, Al, Mn, and Cr, and a part of the carbon element may be substituted with nitrogen atoms. The size and distribution density of the ε-carbide can be observed using a scanning electron microscope, a transmission electron microscope, etc. Mechanical spectrum measurement is useful for detecting small ε-carbide in the initial precipitation stage. The number (density) of ε-carbide with a size of 2 nm or more and 150 nm or less per mm 2 is 1×10 6 or more, preferably 2×10 6 or more, and more preferably 5×10 6 or more. The cooling rate will be described later.

[0033] As a deformed state of the steel material according to the above-described Form 1, it can contain at least one of Nb, Ti, and V, and can contain at least one of Mo and Cr. At least one of Mo and Cr forms a complex carbonitride during the formation of the carbonitride of at least one of Nb, Ti, and V.

[0034] As a deformed state of the steel material according to the above-described Form 1, it can contain (incorporated or absorbed) aluminum nitride in the ε-carbide. Aluminum nitride dissolves and disappears, solid-solves, disperses, or diffuses within the ε-carbide during the growth of the ε-carbide.

[0035] As a deformed state of the steel material according to the above-described Form 1, the mass% ratio of Al to N can be greater than 7. This is to effectively cause slow-diffusing aluminum atoms and fast-diffusing nitrogen atoms to encounter each other, promoting effective nucleation of aluminum nitride and causing it to precipitate in a finely dispersed manner. The mass% ratio of Al to N is preferably 8 or more, and more preferably 10 or more.

[0036] As a deformed state of the steel material according to the above-described Form 1, in terms of volume fraction, the martensite phase or bainite phase can be present in an amount of 70% or more and less than 90%.

[0037] As a deformed state of the steel material according to the above-described Form 1, in terms of volume fraction, the retained austenite phase is 5% or more and less than 30%, and the other phase containing at least the ferrite phase is less than 20%, and it can be present.

[0038] Here, when a steel material having a chemical composition such as that of the steel material according to Mode 1 is rapidly cooled from the austenite phase region in the state diagram at a cooling rate of 0.1 °C / s to 200 °C / s, a steel material having a martensite phase or bainite phase as the main structure, in which aluminum nitride is finely dispersed and precipitated during or after the phase transformation, can be obtained. Thereafter, when tempering is performed in the range of 100 °C to 300 °C, in terms of volume fraction, the martensite phase or bainite phase is 70% or more and less than 90%, the retained austenite phase is 5% or more and less than 30%, and the phase (group of structures) containing at least the ferrite phase and others is less than 20%, and the steel material has a structure represented by these. The volume of each structure is evaluated by electron microscope observation or X-ray diffraction intensity measurement. The volume fraction of the main structure depends on the cooling rate and the local cooling rate of the steel material, but it is necessary to be 70% or more and less than 90% in the martensite phase or bainite phase in order to achieve both high strength and hydrogen embrittlement resistance characteristics.

[0039] As a deformed form of the steel material according to Form 1, the yield strength can be 1000 MPa or more. The yield strength is preferably 1100 MPa or more, and more preferably 1200 MPa or more.

[0040] As a deformed form of the steel material according to Form 1, the tensile strength can be 1200 MPa or more. The tensile strength is preferably 1300 MPa or more, and more preferably 1400 MPa or more.

[0041] As a deformed form of the steel material according to Form 1, the elongation can be 12% or more. The elongation is preferably 13% or more, and more preferably 15% or more.

[0042] As a deformed form of the steel material according to Form 1, a strip-shaped test piece subjected to U-bending is immersed in a 0.1% ammonium thiocyanate solution for at least 100 hours, and it can be such that the end face of the test piece does not break (here, crack fracture). Preferably, it is such that it does not break after being immersed for at least 300 hours, and more preferably, it is such that it does not break after being immersed for at least 500 hours.

[0043] In the present disclosure, the manufacturing method of the steel material according to the following form 2 and its modified forms can be appropriately selected and combined.

[0044] As the manufacturing method of the steel material according to the form 2, A manufacturing method of a steel material for manufacturing the steel material according to claim 1, When cooling from after hot rolling the steel material to room temperature, by cooling so that the martensite phase or bainite phase is formed, aluminum nitride is dispersedly precipitated in the martensite phase or bainite phase; Then, by performing tempering at 100°C or higher and lower than 300°C, the ε-carbide is grown with the aluminum nitride as a nucleus; including, a manufacturing method of the steel material is possible.

[0045] As a modified form of the manufacturing method of the steel material according to the form 2, In the step of growing the ε-carbide, the ε-carbide having a size of 2 nm or more and 150 nm or less is dispersedly precipitated at a density of 1 × 10 2 or more per 1 mm 6 in the martensite phase or bainite phase. This can be achieved.

[0046] As a modified form of the manufacturing method of the steel material according to the form 2, In the step of dispersedly precipitating the aluminum nitride, the cooling rate is 0.1°C / s to 200°C / s. This can be achieved.

[0047] Here, the manufacturing method of the steel sheet according to the form 2 includes a step of dispersedly precipitating aluminum nitride and a step of performing tempering. As a whole, for example, it can be performed in the order of a melting step, a hot rolling step, a cold rolling step, a continuous annealing step, a cooling step (corresponding to the step of forming the martensite phase or bainite phase and dispersedly precipitating aluminum nitride), and a tempering treatment step (corresponding to the step of generating and growing ε-carbide).

[0048] First, in the melting process, for example, a slab is melted from molten steel adjusted to the chemical composition of the steel material according to the above-described Form 1 by a continuous casting method or an ingot-making method.

[0049] In the subsequent hot rolling process, for example, the slab melted in the melting process is once cooled to room temperature, then reheated, and after holding for heat preservation, it is immediately rolled to produce a hot-rolled sheet (hot-rolled steel material). Note that it is also possible to directly roll after casting. When cooling to room temperature and then reheating, the slab heating temperature is 1150°C or higher and it is heated for 1 hour before rolling. The finishing temperature is 950°C or higher. After finishing rolling, it is cooled to approximately 600°C at an average cooling rate of 30°C / s or higher and then coiled.

[0050] In the subsequent cold rolling process, for example, the above hot-rolled sheet is cold-rolled to produce a cold-rolled sheet (cold-rolled steel material) with a predetermined sheet thickness. The reduction ratio is 30% or higher. If the reduction ratio is less than 30%, in the subsequent annealing process, austenite grains may coarsen, and it may not be possible to make the average block diameter of the martensite phase or bainite phase in the steel sheet 5 μm or less.

[0051] In the subsequent continuous annealing process, for example, the obtained cold-rolled sheet is continuously annealed to produce an annealed sheet (annealed steel material). Continuous annealing is preferably performed on a continuous annealing line. In the continuous annealing process, the cold-rolled sheet is heated to a temperature range of Ae3 point - 10°C or higher and 920°C or lower and held for 120 seconds or longer. If the heating and holding temperature is less than Ae3 point - 10°C, the volume fraction of the martensite phase or bainite phase will be less than 70%, and the strength and hydrogen embrittlement resistance characteristics will deteriorate. On the other hand, if the heating and holding temperature exceeds 920°C, austenite grains will coarsen, and it may not be possible to make the average block diameter of the tempered martensite phase or bainite phase 5 μm or less. Note that in the case of a small amount of cold-rolled sheet, batch annealing may also be used.

[0052] In the next cooling process, for example, an annealed plate that has been annealed is cooled to room temperature to 50°C to produce a cooled plate (cooled steel). The cooling rate from the heating and holding temperature to Ms point - 50°C can be 0.1°C / s to 200°C / s. During the cooling to Ms point - 50°C, the transformation to martensite phase or bainite phase and the precipitation of aluminum nitride are approximately completed. The cooling means is not particularly limited and may be any of water cooling, aqueous solution cooling, gas-water cooling, oil cooling, gas cooling, etc. The cooling rate affects the density and size of aluminum nitride precipitated inside the laths of martensite phase or bainite phase and along the dislocation lines. If the cooling rate is too high, the precipitated aluminum nitride is too small to function as the nucleation site of ε-carbide in the tempering process described later. Therefore, the cooling rate is preferably 200°C / s or less, preferably 100°C / s, and preferably 50°C / s. On the other hand, if the cooling rate is less than 0.1°C / s, the transformation to martensite phase or bainite phase is insufficient and an excessive ferrite phase is generated. Therefore, the cooling rate is preferably 0.1°C / s or more, preferably 0.2°C / s or more, and more preferably 0.5°C / s or more. The Ms point of the steel plate can be calculated from the conventionally known Ms point and the empirical formula of the chemical composition, or can also be determined by measuring the thermal expansion curve in the laboratory.

[0053] In the next tempering treatment process, for example, the cooled plate cooled to room temperature to 50°C is inserted into a preheated furnace and held in the temperature range of 100°C or more and less than 300°C for 60 seconds or more and 900 seconds or less to precipitate ε-carbide and obtain a steel plate (steel). If it is less than 60 seconds, the temperature distribution of the steel plate inevitably becomes extremely non-uniform. As a result, the precipitation of ε-carbide becomes non-uniform, and the high strength and hydrogen embrittlement resistance characteristics of the steel plate cannot be ensured. In the case of holding for a long time of 900 seconds or more, the ε-carbide undergoes Ostwald ripening, the interparticle distance of the ε-carbide becomes large, and the precipitation strengthening effect deteriorates. The ε-carbide has a size of 2 nm or more and 150 nm or less, and its distribution density is 1×10 6 pieces / mm 2The above is the case. When the tempering temperature is less than 100°C, the size of ε carbide is less than 2 nm, the effect of precipitation strengthening is small, and the hydrogen embrittlement resistance property is also insufficient. When the tempering temperature is 300°C or higher, even if Si is contained, ε carbide dissolves and disappears in the martensite phase or bainite phase which is the parent phase. Along with this, cementite which is inferior in strength and toughness newly precipitates, and the high strength and excellent hydrogen embrittlement resistance property of the steel sheet are impaired. The tempering temperature is preferably 120°C to 280°C, more preferably 150°C to 250°C.

[0054] The above has described the case where the manufacturing method of the steel material according to the above-described Form 2 is applied to a steel sheet, but it is not limited to the steel sheet and can be applied to steel materials of various shapes such as sections and bars.

[0055] According to the above-described Forms 1 and 2, it is possible to manufacture a steel material excellent in both tensile strength and hydrogen embrittlement resistance property in which delayed fracture due to hydrogen intrusion hardly occurs, and the long-term reliability of the steel material member can be enhanced. Further, according to the above-described Forms 1 and 2, since the amount of steel material used can be reduced, it is possible to reduce the emission of greenhouse gases in the steel refining process and automobile driving, and it can also contribute to the solution of global environmental problems.

Examples

[0056] Hereinafter, examples will be described with reference to the table. Table 1 is a table showing the component composition and Ms point of each steel type (steel material). Table 2 is a table showing the steel type (steel material), heating temperature, holding time, cooling rate, tempering temperature and tempering time, size and density of ε carbide, YS, TS, EL, and delayed fracture resistance property of each sample.

[0057]

Table 1

[0058]

Table 2

[0059] Note that the following examples are merely illustrative and do not limit the present invention. The conditions of the examples are one set of conditions adopted to confirm the feasibility and effects of the present invention. The present invention can adopt various conditions as long as the object of the present invention is achieved without departing from its technical idea and the claims of the invention.

[0060] Steel pieces with a thickness of 30 mm obtained by melting a steel type having the component composition (chemical composition) shown in Table 1 were reheated to 1250 °C and then hot-rolled to a thickness of 3.0 mm at a finishing temperature of 950 °C to produce hot-rolled sheets. After finish rolling the hot-rolled sheets, they were air-cooled to 600 °C and then cooled to a temperature of 100 °C or lower. After pickling the cooled hot-rolled sheets, cold rolling was performed to produce cold-rolled sheets with a thickness of 1.4 mm. Thereafter, annealing, cooling, and tempering treatments were performed on the cold-rolled sheets under the conditions (heating temperature, holding temperature, cooling rate, tempering temperature, tempering time) shown in Table 2 to obtain steel sheets (steel materials). Thereafter, microstructure evaluation, tensile tests, and stress corrosion cracking tests were performed on the steel sheets related to each sample. The test results (YS, TS, EL, stress corrosion cracking resistance characteristics) of the steel sheets related to each sample are shown in Table 2.

[0061] Here, in Table 2, the heating temperature is the heating temperature during annealing, the holding temperature is the holding time during annealing, and the cooling rate is the average cooling rate from the heating and holding temperature to a temperature of Ms point - 50 °C.

[0062] [Tensile Test] From the steel sheets related to each sample, JIS No. 5 test pieces with the major axis in a direction perpendicular to the rolling direction were taken, and tensile tests were performed in accordance with the provisions of JIS Z 2241 (1998). The yield stress (YS: Yield Stress), tensile strength (TS: Tensile Strength), and elongation (EL: Elongation) are shown in Table 2. Here, it is considered to be of high strength when YS ≥ 1000 MPa, TS ≥ 1200 MPa, and EL ≥ 12%.

[0063] [Stress Corrosion Cracking Test] A test piece measuring 100 mm × 30 mm taken with its longitudinal direction parallel to the rolling direction was U-bent with a bending radius of 10 mm, and then stress was applied by tightening the springback portion with bolts (see Non-Patent Document 7). The test piece was immersed in a 0.1% ammonium thiocyanate solution at 25°C, and the time until fracture (here, crack fracture) occurred at the end face portion of the U-bent test piece was investigated to evaluate the stress corrosion cracking resistance characteristics after processing. The 0.1% ammonium thiocyanate solution can introduce hydrogen into the test piece while minimizing the dissolution amount of the test piece during the immersion test. When the test piece did not fracture after being immersed in the 0.1% ammonium thiocyanate solution for 500 hours, the stress corrosion cracking resistance characteristics were considered very good (◎); when it did not fracture after being immersed for 100 hours, the stress corrosion cracking resistance characteristics were considered good (○); and when it fractured, the stress corrosion cracking resistance characteristics were considered poor (×).

[0064] In addition, in the remarks column of Table 2, when the sample had a YS ≥ 1000 MPa, TS ≥ 1200 MPa, EL ≥ 12%, and good or better stress corrosion cracking resistance characteristics, it was described as an example, and the others were described as comparative examples. Also, in the remarks column of Table 1, for the steel types corresponding to each sample in Table 2, those with at least one example were described as examples, and the others were described as comparative examples.

[0065] In order to confirm that the excellent delayed fracture resistance characteristics shown in Table 2 are due to ε-carbide, for the same steel material (Steel type D in Table 1) with the same chemical composition and process up to cooling, an as-annealed sample without annealing treatment and an annealed sample with annealing treatment at 200 °C were prepared. It was confirmed by using an electron microscope that ε-carbide was precipitated in the annealed sample. Both samples were immersed in a 0.1% ammonium thiocyanate solution at 25 °C for 100 hours to introduce hydrogen into the interior of each sample. Then, the hydrogen release rate from each sample was measured while raising the temperature. The results are shown in Fig. 1. In the as-annealed sample (dashed line) where ε-carbide has not yet precipitated, there is a release peak from room temperature to 50 °C to 100 °C. The release peak in this temperature range has conventionally been considered to be due to the desorption of hydrogen atoms trapped at the lath interface, dislocation lines, or metal carbonitride interfaces. In the annealed sample (solid line) where ε-carbide has precipitated, a large release peak appears at 300 °C to 400 °C. This temperature range coincides with the temperature range in which ε-carbide dissolves in the steel structure and cementite is formed. As a result of hydrogen being firmly occluded in ε-carbide, hydrogen is not released to the outside of the sample until the temperature at which the crystal of ε-carbide becomes unstable. The large peak intensity is evidence that a large amount of hydrogen is occluded. In the steel material according to the present invention, the invasive hydrogen is stably occluded inside the crystal of ε-carbide dispersed and precipitated in the martensite phase or bainite phase, so the allowable amount of invasive hydrogen is significantly large. Therefore, the steel material according to the present invention can be stably used in a hydrogen environment up to a temperature of approximately 200 °C.

[0066] In this application, it has been shown that by dispersing and precipitating ε-carbide composed of iron and carbon, which are the basic elements of steel, in a steel material, it is possible to produce a steel material with high strength and excellent hydrogen embrittlement resistance without using expensive rare elements. Although ε-carbide in steel was discovered in the 1940s, it was substantially ignored as a carbide that only transiently exists in the phase diagram and was only mentioned at the end of a group of iron carbides as a superordinate concept in domestic and foreign patent documents. Even if it was rarely the subject of basic research, there was no exploration of its technical and industrial value. However, in the early 21st century, a military steel material called "Eglin Steel" with high strength and high toughness, which requires less added alloy, was developed at the United States Air Force Headquarters in Eglin, Florida, USA, as the warhead material for a penetrator bomb. Subsequently, it was revealed by electron microscope observation that the dispersed and precipitated ε-carbide contributes to high strength and high toughness near room temperature. The fact that this steel has become the focus of interest along with ε-carbide in some American academic societies was finally learned by the inventor of this application during the correspondence with the referee when submitting the manuscript of Non-Patent Document 1 to an American journal. From this vein, the general-purpose special steel with high strength and high toughness by dispersed ε-carbide listed in Patent Document 3 was born. However, its production has the drawback that it requires a special melting furnace and complex heat treatment. The invention of this application is based on a novel and original idea and attempts to produce a general-purpose steel with high strength (and high toughness) and high hydrogen embrittlement resistance by making full use of the still advanced steel technology in our country. Its technical spillover effect is significant.

[0067] After filing the Japanese application of the present application, the original paper of Non-Patent Document 2 by T.G.O. Berg described in

[0014] (Non-Patent Document 8 by T.G.O. Berg) could be obtained. The Berg original paper was not available at the time of preparing the Japanese application of the present application because the library was locked out due to the spread of COVID-19. Non-Patent Document 8 (oral presentation in 1959, and later published as a paper in 1961) is the mother paper of Non-Patent Document 2 (published in 1962). As a result of a detailed examination of Non-Patent Document 8, it was found that Berg's experiment was very incomplete as a crystal chemistry experiment. Furthermore, in 1973, it was finally determined by experts in X-ray diffraction that the carbide assumed by Berg was not the hexagonal ε-carbide, but Fe5C2, a carbide with a different crystal structure. In these two respects, Berg did not discuss the ε-carbide. These circumstances have been recently published as a research paper by the inventor of the present application (published on February 23, 2021, Reference 1). The entire content thereof shall be incorporated herein by reference. (Reference 1) Michio Shimotomai, "Heuristic Design of Advanced Martensitic Steels That Are Highly Resistant to Hydrogen Embrittlement by ε-Carbide" Metals, 2021, 11(2):370. https: / / doi.org / 10.3390 / met11020370

[0068] Incidentally, the disclosures of the above-mentioned patent documents and non-patent documents are hereby incorporated by reference into this document and can be used as the basis or part of the present invention as necessary. Within the framework of the entire disclosure of the present invention (including the claims and the drawings), further modifications and adjustments of the embodiments or examples can be made based on its basic technical concept. Also, within the framework of the entire disclosure of the present invention, various combinations or selections (including non-selections if necessary) of various disclosure elements (including each element of each claim, each element of each embodiment or example, each element of each drawing, etc.) are possible. That is to say, the present invention naturally includes all disclosures including the claims and the drawings, and various deformations and modifications that could be made by those skilled in the art in accordance with the technical concept. Further, regarding the numerical values and numerical ranges described in this application, even if not explicitly stated, any intermediate value, lower-order numerical value, and small range are considered to be described. Furthermore, each disclosed matter of the above-cited documents is, as necessary, considered to be included in (belong to) the disclosed matters of this application as part of the disclosure of the present invention of this application, and can be used in combination with the matters described in this document, either in part or in whole, in accordance with the spirit of the present invention of this application.

Claims

1. By mass percentage, C: 0.15% to 0.35%, Si: 0.8% to 2.5%, Mn: 0.8% to 2.5%, Al: 0.03% to 2.0%, N: 0.002% to 0.010%, P: 0.01% or less, S: 0.01% or less, O: 0.01% or less, B: 0.0001% to 0.005%, Nb: 0.0% to 0.05%, Ti: 0.0% to 0.2%, V: 0.0% to 0.05%, Mo: 0.0% to 1.0%, Cr: 0.0% to 1.0%, Ni: 0.01% to 1.0%, Cu: 0.05% to 1.0%, At least one of Ca, Mg and REM (rare earth metals): 0.0005% to 0.01%, and, the balance: Fe and impurities, having a chemical composition represented by, The martensite phase and / or bainite phase in which ε carbide having a size of 2 nm or more and 150 nm or less is dispersed and precipitated at a density of 1 × 10 or more per 1 mm ２ per ６ has by volume fraction, the martensite phase and / or bainite phase being present in an amount of 70% or more and less than 90%, a steel material, the ε-carbide having a hexagonal crystal structure containing carbon vacancy sites that capture hydrogen within its lattice structure, the carbon vacancy sites that capture hydrogen being stable from 0°C to about 300°C, and the steel material substantially not containing cementite, a steel material characterized by the above.

2. containing more than 0 mass% of at least one of the Nb, the Ti and the V, containing more than 0 mass% of at least one of the Mo and the Cr, the steel material according to Claim 1.

3. the ratio by mass percentage of the Al to the N being greater than 7, or 10 or more, the steel material according to Claim 1 or 2.

4. by volume fraction, the retained austenite phase being 5% or more and less than 30%, and the other phase containing at least ferrite being less than 20%, present, the steel material according to any one of Claims 1 to 3.

5. the steel material having a hydrogen heat release rate characteristic curve having a larger second peak in the temperature range of 300°C to 400°C compared to the first peak that appears per 100°C in the hydrogen heat release rate characteristic curve in the range of 0°C to 400°C, the second peak being generated by either or both of the release of hydrogen stably captured by the hydrogen capture carbon vacancy sites in the ε-carbide and the transformation of the ε-carbide, the steel material according to any one of Claims 1 to 4.

6. the ε-carbide having substantially no hydrogen capture ability after heating at 400°C, as far as detectable by the hydrogen heat release rate characteristic curve, the steel material according to any one of Claims 1 to 5.

7. The steel material according to any one of claims 1 to 6, which substantially does not contain cementite in a temperature range of 0°C to about 300°C.

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